Compositions: ceramic – Ceramic compositions – Pore-forming
Reexamination Certificate
2002-05-30
2004-12-21
Silverman, Stanley S. (Department: 1754)
Compositions: ceramic
Ceramic compositions
Pore-forming
Reexamination Certificate
active
06833337
ABSTRACT:
BACKGROUND
Refractory ceramics (e.g., aluminates, aluminosilicates) can exhibit several enhanced properties relative to refractory metals and alloys, such as corrosion resistance, high temperature stability in oxidizing atmospheres, specific strength and stiffness, creep resistance, and wear resistance. However, the brittleness of ceramic bodies renders the fabrication of shaped components tedious and expensive. Such brittleness also necessitates the use of reinforcements to enhance damage tolerance.
Despite extensive research and development over the past few decades, continuous-ceramic-fiber-reinforced, ceramic-matrix composites have not found widespread use in high-temperature structural applications involving oxidizing atmospheres (e.g., turbine engine applications). The use of these materials has been hampered by the inability to produce low-cost, net-shaped composites that are capable of retaining both a high strength (i.e., owing to a high fiber strength) and a high toughness (i.e., owing to sufficiently weak fiber-matrix interface for fiber-pullout toughening) under such high-temperature, oxidizing conditions (e.g., ≈1200° C. in high-pressure air for turbine engine applications) [
1
-
7
]. Erosion/oxidation resistance are of particular concern in the very high-temperature (≈2000° C.) corrosive conditions of liquid-fueled rocket engines [
7
]. Although continued research and development of high-strength oxide fibers and oxidation-resistant interface coatings may ultimately enable the low-cost fabrication of oxide fiber-reinforced ceramic composites with satisfactory damage tolerance and erosion/oxidation resistance for jet and rocket engine applications, oxide-matrix composites with other types of reinforcements should be considered in the interim.
An alternate approach for reinforcing oxide matrices is to use metallic (or intermetallic) alloys that are oxidation resistant, such as Ni-based compositions, or very high-melting, such as Nb-based compositions (Nb, Nb—Al solid solutions, and Nb
3
Al melt at 2468° C., 2060-2468° C., and 1940-2060° C., respectively [
8
,
9
]). As reinforcement materials, Ni-rich or Nb-rich solid solutions, intermetallic compounds, or mixtures thereof offer a number of attractive features.
For example, owing to their oxidation resistance at elevated temperatures, damage tolerance (e.g., yield strength, ultimate tensile strength, toughness, fatigue resistance) over a range of temperatures, and capability for being fabricated into complex shapes (e.g., by casting), Ni-based superalloys are used extensively for high-temperature components in engine applications (e.g., combustor liners and ducts, disks and blades, and nozzle vanes in turbine engines; exhaust valves in reciprocating engines; pre-combustion chambers in diesel engines) [
5
-
7
,
10
-
13
]. Aluminum additions are used to enhance the high-temperature strength of such superalloys, owing to the formation of the ordered compound □′-Ni
3
Al as a coherent precipitate dispersed within the Ni-rich (FCC □ phase) solid solution matrix [
1
,
5
,
11
,
14
-
19
]. Pure □′-Ni
3
Al and its alloys exhibit higher yield strengths at elevated temperatures than the □ phase [
11
,
14
-
19
]. The □′ compound can also act as a ductile reinforcement material in composites (at room temperature or elevated temperatures), either as a polycrystalline material with proper alloying (e.g., Al concentration <25 at % with dopants such as boron and chromium) or in single crystal form (e.g., as single crystal filaments) [
11
,
14
-
18
]. Stoichiometric □-NiAl, while less ductile at room temperature than Ni-based superalloys and □′-Ni
3
Al alloys, possesses a higher elastic modulus, a higher thermal conductivity, and a higher melting point (1638° C.) [
15
,
19
-
23
]. Niobium alloys and intermetallic compounds (e.g., Nb
3
Al) are very high melting and exhibit better creep resistance than Ni
3
Al, possess high thermal conductivities at elevated temperatures, and are less dense than other refractory metals (such as W, Ta, Mo) [
8
,
24
-
26
].
As indicated by the discussion above, these Ni-bearing and Nb-bearing alloys offer a tailorable range of thermo-mechanical properties that can be exploited in ceramic composites. In addition, the negative features of these alloys (i.e., weight, creep) can be significantly reduced if such alloys are used as reinforcements within composites containing a stiff, continuous, light-weight ceramic phase.
Accordingly, there remains a need for a method for the fabrication of near net-shaped AEAl
2
O
4
-bearing (where AE=Mg, Ba) composites reinforced with M—Al (M=Ni, Nb) solid solutions and/or intermetallic compounds (e.g., Nb, NiAl+Ni
3
Al, or Ni
3
Al+Ni-based solid solutions) [
27
].
Because the ionic species in these aluminates have essentially fixed valence states, these oxide compounds are thermodynamically compatible with oxygen at high temperatures. On the other hand, Nb, Ni, (Nb,Al) or (Ni,Al) solid solutions, and □′-Ni
3
Al can act as ductile reinforcements at ambient and elevated temperatures (□-NiAl also becomes ductile above ~400° C.) [
15
,
17
,
19
,
23
]. The thermal conductivities of Nb, Ni, and Ni—Al alloys (particularly □-NiAl) are much higher than for the aluminates [
25
]. The coefficients of thermal expansion (CTE) of the aluminates are smaller than for nickel or the nickel aluminides, which should place the ceramic phase in compression upon cooling from the peak processing or use temperature [
9
,
32
]. For very high-temperature applications, composites of AE aluminate and Nb or Nb
3
Al are attractive in that these phases possess similar CTE values (8.2-9.2×10
−6
/° C. [
9
,
32
]).
Although the properties of AEAl
2
O
4
/M—Al alloy composites will depend on factors such as the amounts, sizes, and distributions of phases, and the degree of interfacial bonding, the discussion above (and in the following sections) indicates that such composites can be:
lighter and more resistant to high-temperature creep, oxidation, and erosion relative to monolithic metallic alloys or intermetallics,
more fracture resistant (tougher, higher fracture strengths) and thermally conductive than monolithic aluminates,
formed into near net shapes by the process of the present invention as described herein.
Accordingly, and with the foregoing objectives and advantages in mind, the present invention is summarized below. In view of the present disclosure or through practice of the present invention, other advantages may become apparent.
SUMMARY OF THE INVENTION
A variety of reaction-based techniques have been developed for the in-situ syntheses of monolithic ceramics and ceramic composites, including the Co-Continuous Ceramic Composite (C
4
) process, the reactive metal penetration process (RMP) [
3
-
5
], the Directed Metal Oxidation (DIMOX) [
6
,
7
], the Reaction Bonded Metal Oxide (RBMO) process [
8
,
9
], Self-Propagating High-Temperature Syntheses (SHS) [
10
], Hillig or Hucke Process [
11
-
13
], and the Solid Metal-Bearing Precursor (SMP) process [
14
-
18
].
Such AEAl
2
O
4
/M—Al alloy composites should exhibit enhanced thermo-mechanical properties relative to monolithic AEAl
2
O
4
or monolithic M—Al alloy bodies. Compared to Nb, Ni, and Ni—Al alloys, polycrystalline MgAl
2
O
4
and BaAl
2
O
4
possess higher values of elastic moduli, lower densities, and, hence, higher values of specific stiffness (e.g., E/□=79, 40, 24, 23, and 12 MPa·cm
3
/g for MgAl
2
O
4
, NiAl, Ni
3
Al, Ni, and Nb respectively, at room temperature) [
11
,
15
,
22
,
23
,
28
]. These stiff, stoichiometric aluminates exhibit better creep resistance at elevated temperatures than Nb, Ni, or Ni—Al-based solid solutions or intermetallic compounds [
29
-
31
&rsqb
Kumar Pragati
Sandhage Kenneth H.
Johnson Edward M.
Silverman Stanley S.
Standley Law Group LLP
The Ohio State University
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